Composites having an intermetallic containing matrix

ABSTRACT

This invention relates to a composite material comprising an in-situ precipitated second phase in an intermetallic matrix, and to the process for making such a composite.

This application is a Continuation-In-Part of U.S. patent applicationSer. No. 662,928, filed Oct. 19, 1984 now abandoned.

The present invention relates generally to a process for formingcomposite materials and to a composite product having an intermetalliccontaining matrix including an in-situ precipitation of a second phase,such as another intermetallic phase or a ceramic material, wherein thesecond phase comprises a boride, carbide, oxide, nitride, silicide,sulfide, etc., or intermetallic of one or more metals.

BACKGROUND OF THE INVENTION

For the past several years, extensive research has been devoted to thedevelopment of metal-ceramic composites, such as aluminum reinforcedwith carbon, boron, silicon carbide, silica, or alumina fibers,whiskers, or particles. Metal-ceramic composites with good hightemperature yield strengths and creep resistance have been fabricated bythe dispersion of very fine (less than 0.1 micron) oxide or carbideparticles throughout the metal or alloy matrix. However, this metalceramic composite technology has not heretofore been extended to includeintermetallic matrices. Prior art techniques for the production ofmetal-ceramic composites may be broadly categorized as powdermetallurgical approaches, molten metal techniques, and internaloxidation processes.

The powder metallurgical type production of such dispersion-strengthenedcomposites would ideally be accomplished by mechanically mixing metalpowders of approximately 5 micron diameter or less with the oxide orcarbide powder (preferably 0.01 micron to 0.1 micron). High speedblending techniques or conventional procedures such as ball milling maybe used to mix the powder. Standard powder metallurgy techniques arethen employed to form the final composite. Conventionally, however, theceramic component is large, i.e., greater than 1 micron, due to a lackof availability, and high cost, of very small particle size materialssince their production is energy intensive, time consuming, and costlyin capital equipment. Furthermore, production of very small particlesinevitably leads to contamination of the particles with oxides,nitrides, and materials from various sources such as the attritor (e.g.,iron). The presence of these contaminants inhibits particulate-to-metalbonding which in turn compromises the mechanical properties of theresultant composites. Further, in many cases where the particulatematerials are available in the desired size, they are extremelyhazardous due to their pyrophoric nature.

Alternatively, it is known that proprietary processes exist for thedirect addition of appropriately coated ceramics to molten metals.Further, molten metal infiltration of a continuous ceramic skeleton hasbeen used to produce composites. In most cases, elaborate particlecoating techniques have been developed to protect the ceramic particlesfrom the molten metal during admixture or molten metal infiltration, andto improve bonding between the metal and ceramic. Techniques such asthese have resulted in the formation of silicon carbide-aluminumcomposites, frequently referred to as SiC/Al, or SiC aluminum. Thisapproach is only suitable for large particulate ceramics (e.g., greaterthan 1 micron) and whiskers, because of the high pressures involved forinfiltration. The ceramic material, such as silicon carbide, is pressedto form a compact, and liquid metal is forced into the packed bed tofill the intersticies. Such a technique is illustrated in U.S. Pat. No.4,444,603, of Yamatsuta et al, issued Apr. 24, 1984. Because of thenecessity for coating techniques and molten metal handling equipmentcapable of generating extremely high pressures, molten metalinfiltration has not been a practical process for making metal-ceramiccomposites.

The presence of oxygen in ball-milled powders used in prior art powdermetallurgy techniques, or in molten metal infiltration, can result inoxide formation at the interface between the ceramic and the metal. Thepresence of such oxides will inhibit interfacial binding between theceramic phase and the matrix, thus adversely effecting ductility of thecomposite. Such weakened interfacial contact can also result in reducedstrength, loss of elongation, and facilitated crack propagation. Inaddition, the matrix may be adversely effected, as in the case oftitanium which is embrittled by interstitial oxygen.

Because of the above-noted difficulties with conventional processes, thepreparation of metal-ceramic composites with submicron ceramicdispersoids for commercial applications has been extremely expensive.

Internal oxidation of a metal containing a more reactive component hasalso been used to produce dispersion strengthened metals, such asinternally oxidized aluminum in copper. For example, when a copper alloycontaining about 3 percent aluminum is placed in an oxidizingatmosphere, oxygen may diffuse through the copper matrix to react withthe aluminum, precipitating alumina. This technique, although limited torelatively few systems since the two metals utilized must have a widedifference in chemical reactivity, has offered a feasible method fordispersion hardening. However, the highest possible level of dispersoidsformed in the resultant dispersion strengthened metal is generallyinsufficient to impart significant changes in properties such asmodulus, hardness, and the like. In addition, oxides are typically notwetted by the metal matrix, so that interfacial bonding is not optimum.

In recent years, numerous ceramics have been formed using a processreferred to as self-propagating high-temperature synthesis (SHS), whichinvolves an exothermic, self-sustaining reaction which propagatesthrough a mixture of compressed powders. Generally, the SHS process isignited by electrical impulse, thermite, or spark. The SHS processinvolves mixirg and compacting powders of the constituent elements, andigniting the green compact with a suitable heat source. On ignition,sufficient heat is released to support a self-sustaining reaction, whichpermits the use of sudden, low power initiation of high temperatures,rather than bulk heating over long times at lower temperatures.Exemplary of these techniques are the patents of Merzhanov et al. InU.S. Pat. No. 3,726,643, there is taught a method for producinghigh-melting refractory inorganic compound by mixing at least one metalselected from groups IV, V, and VI of the Periodic System with anon-metal such as carbon, boron, silicon, sulfur, or liquid nitrogen,and locally heating the surface of the mixture to produce a localtemperature adequate to initiate a combustion process. In U.S. Pat. No.4,161,512, a process is taught for preparing titanium carbide bylocalized ignition of a mixture consisting of 80-88 percent titanium and20-12 percent carbon, resulting in an exothermic reaction of the mixtureunder conditions of layer-by-layer combustion. These references dealwith the preparation of ceramic materials, in the absence of a secondnon-reactive metallic phase.

U.S. Pat. No. 4,431,448 teaches preparation of a 35rd alloy byintermixing powders of titanium, boron, carbon, and a Group I-B bindermetal, such as copper or silver, compression of the mixture, localignition thereof to initiate the exothermic reaction of titanium withboron and carbon, and propagation of the reaction, resulting in an alloycomprising titanium diboride, titanium carbide, and the binder metal.This reference, however, is limited to the use of Group I-B metals suchas copper and silver, as binders. As is set forth in the patent,products made by this method have low density, requiring subsequentcompression and compaction.

Another class of materials which has seen considerable interest anddevelopment is intermetallic materials, especially intermetallics ofaluminum such as the aluminides of titanium, zirconium, iron, cobalt,and nickel.

The need for the advanced properties obtainable with intermetallicmaterials is typified by their potential application to structurescapable of withstanding high temperatures, such as turbine engines. Indesigning and operating turbine engines today and for the foreseeablefuture, there are two primary problem which demand solutions from thefield of materials science. The first of these is the need to operatecertain portions of the engine at higher gas and metal temperatures toimprove operating efficiency and save fuel. The second problem is theneed for lighter materials to decrease engine weight and engineoperating stresses due to heavy rotating components, and to increase theoperating life of disks, shafts, and bearing support structures. Theselatter structures require materials which are less dense than the nickelbase superalloys they are intended to replace, but which possess roughlythe same mechanical properties and oxidation resistance as thosematerialsin current usage.

The intermetallics are typically highly ordered compounds, in the sensethat they possess regularly repeating (e.g., A B A B A B) atomsequencing. Intermetallic compounds are particularly suited to theseneeds because of two properties which derive from the fact that theypossess ordered structures. Modulus retention at elevated temperature inthese materials is particularly high because of strong A-B bonding. Inaddition, a number of high temperature properties which depend ondiffusive mechanisms, such as creep, are improved because of thegenerally high activation energy required for self-diffusion in orderedalloys.

The formation of long range order in alloy systems also frequentlyproduces a significant positive effect on mechanical properties,including elastic constants, strength, strain-hardening rates, andresistance to cyclic creep deformation. Finally, in the case ofaluminides, the resistance to surface oxidation is particularly goodbecause these materials contain a large reservoir of aluminum that ispreferentially oxidized.

However, during metallurgical processing, one problem encountered isthat these materials tend to form coarse grains, which degrade certainmechanical properties, the most important of which is ductility. Also,in many intermetallics the strong A-B bonding results in low temperaturebrittleness, although the exact mechanism of the ductile-brittletransition seems to be different for the different intermetalliccompounds. It is thus necessary to address the problem of minimal lowtemperature ductility without destroying the inherent high temperaturestrength and stiffness. In the prior art it has generally beenconsidered that these latter high temperature properties may only beretained by preserving the ordered structure, hence sacrificing lowtemperature ductility.

Since the early 1970's, the pace of work on ordered alloys andintermetallic compounds has slackened, as a result of lack of progressin improving either ductility or creep resistance of these otherwisevery intriguing alloys.

Interest in utilizing ordered alloys for structural applications wasreawakened in this country when researchers discovered that ductilityand strength improvements could be achieved in TiAl and Ti₃ Al basedalloys using a combination of powder metallurgy and alloying techniques.Later work on the titanium aluminides utilized ingot metallurgy. Thedevelopment of rapid solidification methods led to renewed interest inthe iron and nickel aluminides. The replacement of cobalt in Co₃ V bynickel, and then iron, led to a series of face-centered cubic Ll₂ -typesuperlattices with greater ductility at ambient temperatures. Also, ithas been reported in Japan that polycrystalline Ni₃ Al can be made moreductile by adding small quantities of boron. Later, this work wasconfirmed and the critical composition range over which boron wasbeneficial was identified. (See U.S. Pat. No. 4,478,791 of Huang et al,assigned to General Electric.) These discoveries, together with thenational search for replacements for strategic metals, such as cobaltand chromium, and the need to develop energy-efficient systems, have inthe past few years or two stimulated much additional work; largely inthe area of improving low temperature ductility and increasing hightemperature strength.

Despite these efforts, little progress has been made in developingpractical intermetallic compositions that have sufficiently improved lowtemperature ductility while maintaining high temperature strength.

SUMMARY OF THE INVENTION

It is an object of the present invention to provide a method for formingcomposite materials of discretely dispersed particulate second phasematerials in intermetallic containing matrices, particularly inaluminide containing matrices. The dispersed material may constitute asecond phase such as a ceramic, or an intermetallic compound other thanthe matrix material.

It is a further object of this invention to provide a method fordispersion strengthening of intermetallics such as aluminides. It is aparticular object of this invention to provide a method for theformation of one or more nitride, boride, sulfide, silicide, oxide, andcarbide particulates in a matrix of one or more intermetallic materials.

It is yet a further object of the invention to produced compositeshaving an intermetallic containing matrix which has fine grains forimproved ductility and mechanical properties while retaining the hightemperature characteristics of the intermetallics.

It is also an object of the present invention to provide anintermetallic composite material which may be subjected to conventionalmetallurgical processing steps, such as remelting, annealing, working,extrusion, etc.

Generally, the present invention relates to a process for formingcomposite materials comprising finely divided ceramic or other secondphase particles in an intermetallic containing matrix by an in-situprecipitation of up to about 95 percent by volume of ceramic material inthe matrix or precursors thereof, wherein the ceramic comprises aboride, carbide, oxide, nitride, oxynitride, silicide, sulfide,oxysulfide or a mixture thereof. It has been found that by mixing theconstituents or elements of the desired second phase material with asolvent matrix material comprising an intermetallic or precursorsthereof, and heating to a temperature at which an exothermic reactionwhich forms the second phase is initiated, a solvent assisted reactionensues, resulting in the extremely rapid formation and dispersion offinely divided particles of the second phase material in the matrixmaterial. Where the reaction takes place in a single metal which is aprecursor of an intermetallic, a subsequent reaction or dilution isrequired to convert the matrix material to the intermetallic.

The invention further relates to a process for forming compositematerials comprising one or more second phase materials in anintermetallic containing matrix material, such as an aluminide, byproviding a substantially molten mass containing the intermetallic orprecursors thereof and then adding at least one of the constituents orelements of the desired second phase ceramic material to the moltenmass, thereby initiating the solvent assisted in-situ precipitationreaction, to form and disperse finely divided particles of the secondphase material in the matrix material.

The invention also relates to a process for forming intermetallic matrixcomposite materials comprising precipitating at least one second phasematerial by contacting reactive second phase forming constituents, inthe presence of a solvent matrix material comprising an intermetallic orprecursors thereof in which said constituents are more soluble than saidsecond phase, at a temperature at which sufficient diffusion of saidconstituents into said solvent matrix material occurs to initiate thereaction of said constituents to produce a material comprising finelydivided particles of the second phase material in a matrix materialcontaining an intermetallic or a precursor thereof, and then introducingthe thus produced composite material into either a molten metal, or amolten intermetallic containing material, wherein said molten metal isat least partially converted to an intermetallic compound or a mixtureof intermetallic compounds.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 represents a schematic presentation of some of the variousreaction modes and states which may be used with this invention.

FIG. 2 is a photomicrograph of a dispersion of titanium diboride in atitanium aluminide (Al₃ Ti) matrix.

FIG. 3 is a photomicrograph of a dispersion of titanium diboride in atitanium aluminide (Al₃ Ti) matrix after etching of the aluminidematrix.

FIG. 4 is a photomicrograph of a dispersion of titanium diboride intitanium aluminide (Ti₃ Al) matrix after etching of the aluminidematrix.

FIG. 5 is a photomicrograph of a dispersion of titanium diboride intitanium aluminide (TiAl) after etching of the aluminide matrix.

FIG. 6 is a photomicrograph of a dispersion of titanium diboride innickel aluminide after etching of the aluminide matrix.

DESCRIPTION OF PREFERRED EMBODIMENTS

As was mentioned previously, the present invention relates to a processfor the in-situ precipitation of fine particulate ceramics or othersecond phases, such as refractory hard metal borides and intermetallicdispersoids within intermetallic systems or precursors thereof, torecover an intermetallic-second phase composite having enhancedmechanical properties, such as high elastic modulus, high-temperaturestability, ductility, and improved wear resistance. However, the processdescribed may also be employed for producing larger particles of thesecond phase material in the chosen intermetallic containing matrix, upto the point at which such larger particles result in componentembrittlement, or loss of ductility, etc. The enhanced mechanicalproperties offer weight-savings in stiffness limited applications,higher operating temperatures and associated energy efficiencyimprovements, and reduced wear in parts subject to erosion. Ofparticular importance in the case of intermetallics is the property offine grain size which imparts greater ductility to these materials thanheretofore attainable with intermetallics processed by prior arttechniques. While the grain size of the intermetallic matrix ofcomposites of the present invention may be from 0.01 to 10 microns orlarger, the preferred grain size range is from 0.01 to 5 microns, andthe most preferred range is from 0.01 to 1 micron. A specificapplication of such materials is in the construction of structuralcomponents capable of high temperature service, such as turbine blades.

Broadly stated, the present invention constitutes a process wherebyelements forming a second phase precipitate react in a solvent matrixmaterial containing an intermetallic, or at least one precursor thereof,to form a finely-divided dispersion of the second phase material in thesolvent matrix material. In the case of aluminum-ceramic reinforcedcomposites, it has been found that the reaction commences, or isinitiated, at a temperature far below the temperature conventionallyrequired for the reaction sought absent the solvent matrix material.While ceramic materials constitute the preferred second phase fordispersion as a fine precipitate in the intermetallic matrix, it is alsopossible to precipitate an intermetallic material as the second phasedispersoid in the intermetallic matrix. In such instances, theprecipitate and matrix must be of different intermetallic composition.While the discussion herein shall focus upon ceramic materials as thesecond phase, or dispersoid, it must be borne in mind that intermetallicsecond phases are also to be included in the scope of the presentinvention. The second phase-forming constituents most easily combine ator about the melting temperature of the solvent matrix material, and theexothermic nature of this reaction may cause a very rapid temperatureelevation or spike, which has the effect of melting additional matrixmaterial, simultaneously causing the further reaction of the secondphase-forming constituents.

In practicing this process, particularly for modulus limitedapplications, it is generally advisable to produce a composite materialcontaining at least about 10 volume percent, preferably about 15 volumepercent, second phase material, so as to yield a composite havingappreciably improved properties in this regard. If grain refining aloneis desired, lower levels of second phase material are adequate. Bestresults appear to be achieved when the concentration of the second phasein the produced composite material is great enough that theinterparticle spacing is one micron or less. The composite may compriseup to about 95 volume percent dispersoid, where further dilution ofcomposite by the addition thereof to an intermetallic or other metalmelt is contemplated. In general, the composite may comprise from about0.5 volume percent to about 25 volume percent dispersoid, with fromabout 1 percent to about 10 volume percent dispersoid preferred forcontrol of matrix grain size.

Exemplary of suitable second phase precipitates are the borides,carbides, oxides, nitrides, silicides, oxynitrides, sulfides, andoxysulfides. Suitable elements include all of the elements which arereactive to form ceramic precipitates, including, but not limited to,transition elements of the third to sixth groups of the Periodic Table.Particularly useful ceramic phase forming constituents include aluminum,titanium, silicon, boron, carbon, oxygen, nitrogen, sulfur, molybdenum,tungsten, niobium, vanadium, zirconium, chromium, hafnium, cobalt,nickel, iron, magnesium, tantalum, manganese, zinc, lithium, beryllium,thorium, and the rare earth elements including scandium, yttrium,lanthanum and the lanthanide series elements such as cerium and erbium.Reactive compounds of such elements, such as B₂ O₃, and B₄ C, and BN mayalso be used.

It is especially to be noted that plural dispersoids, and/or complexcompounds such as titanium zirconium boride, may advantageously beprecipitated in-situ in the intermetallic matrix. For example,composites of relatively low oxygen content may be produced byincorporation of small amounts (e.g., up to about 5 percent, dependentupon oxygen content of the matrix) of strong oxide formers, such asyttrium or any of the rare-earth metals, e.g., cerium and erbium, whichwill scavenge oxygen from the solvent matrix material. The exothermictemperature rise of the reaction mass, in conjunction with the increasedsurface area of the dispersoid formed, may effectively improvescavenging capability. The oxide particles thus formed enhance matrixductility by removal of interstitial oxygen, and may also serve todispersion strengthen the matrix and reduce grain size therein, in turnfurther enhancing matrix ductility. Further, it is to be noted that inmany intermetallic-ceramic composites prepared by the method of thepresent invention, intermetallic precursors will also react with ceramicconstituents to form additional ceramic dispersoids.

As the matrix or solvent, one may use any intermetallic, or precursorsthereof capable of dissolving or at least sparingly dissolving theconstituents forming the ceramic or other second phase, and having alesser capability for dissolving or otherwise reacting with the formedceramic or other second phase precipitate. Thus, at the temperaturesexperienced during the process, the matrix component must act as asolvent for the second phase reaction species, but not for the desiredsecond phase precipitate. It is especially to be noted that the initialmatrix acts primarily as a solvent in the process, and that theconstituents of the second phase precipitate have a greater affinity foreach other than either has for the solvent matrix material. It isfurther to be noted that the intermetallic precursors may individually,or collectively, act as solvent materials for the ceramic or secondphase precursor materials. Thus, the ceramic or other second phaseprecursors, but not the ceramic or second phase itself, must be solublein at least one of the intermetallic precursors or the intermetalliccompound itself. Additionally, it is important that the exothermici ofthe second phase forming reaction be sufficient to sustain the reactionof the second phase precursors and allow it to go to substantialcompletion by consuming the reactants. Therefore, while the potentialchoice of second phase dispersoids and matrix materials is large, thischoice is limited by adherence to the criteria hereinabove recited.

The solvent metal precursors for the intermetallic matrix may beselected from the group consisting of aluminum, nickel, copper,titanium, cobalt, iron, platinum, gold, silver, niobium, tantalum,boron, zinc, molybdenum, yttrium, hafnium, tin, tungsten, lithium,magnesium, beryllium, thorium, silicon, chromium, vanadium, zirconium,manganese, scandium, lanthanum, and rare earth elements and alloysthereof. Preferred intermetallic precursors include aluminum, nickel,titanium, cobalt, iron, and refractory metals. Plural intermetallicmaterials may, of course, be present in the matrix. It is noted that theterms intermetallic containing matrix, or intermetallic matrix, as usedherein, are meant to define a matrix which is predominantlyintermetallic, although other materials, e.g., metal intermetallicprecursors, may also be present in lesser amount.

Intermetallics are suitable in the present invention for both thematrix, and optionally the second phase dispersoid. In such instance,the intermetallic dispersoid may be prepared from the preceding group ofintermetallic matrix precursors, with the exception of copper, gold,silver, and platinum. These materials are generally considered ascompounds of two or more metals in substantially stoichiometricproportions which do not necessarily follow normal ionic/covalentbonding type valence rules.

Intermetallics generally can be defined as phases or compounds of thegeneral formula A_(x) B_(y), wherein A is a metallic element, B isanother metallic element (aluminum in the case of aluminides) and x andy closely approximate integers indicating that the compounds are acombination of two metals in defined molecular proportions. Among themetallic elements capable of forming aluminides are titanium, nickel,iron, cobalt, and refractory metals such as niobium, zirconium, tantalumand the like. Titanium forms the aluminides Ti₃ Al, TiAl, and Al₃ Ti,while nickel forms Ni₃ Al and NiAl. Other aluminides include Zr₃ Al, Co₃Al and Fe₃ Al. For the purposes of the present invention, the aluminidesof titanium and nickel are presently preferred. Substitution of one ormore elements within an intermetallic compound is possible, and may bedesirable to effect crystal lattice changes. Exemplary is thesubstitution of aluminum by titanium in Ni₃ Al to form Ni₃ (Al, Ti). Inaddition, two phase mixtures such as TiAl and Ti₃ Al are suitable.

It should be recognized that aluminides are not necessarily of acomposition such that x and y of the above formula are exact integers.For example, nickel aluminide is commonly referred to as Ni₃ Al althoughit is an intermetallic phase and not a simple ionic or covalently bondedcompound, as it exists over a range of compositions as a function oftemperature, e.g., about 72.5 to 77 weight percent nickel (85.1 to 87.8atomic percent) at about 600° C. Thus, aluminides, and intermetallicsgenerally, may be defined as the compounds which form uponsolidification of near stoichiometric amounts of the appropriate metals.In the molten state, however, the intermetallics exist primarily as arandom mixture of the elements thereof, possessing only relatively shortrange order. Within the scope of the present invention, this randommixture, or intermetallic derived liquid, may act as a solvent matrixmaterial through which the solvent assisted reaction of second phaseprecursors occurs. This molten state may thus be referred to as an"intermetallic derived solvent", or "solvent matrix material" whichterms also encompass the molten state of one or more precursors of saidintermetallic.

With reference to composites with an aluminide containing matrix made bythe methods of the present invention, such composites tend to have ahigher ductility than even those aluminides produced by state-of-the-artprocessing routes absent dispersoids. The aluminide composites are alsocharacterized by fine grain size, which is thought to increase theambient temperature ductility by reducing dislocation stress at grainboundaries as a result of reduced slip length. The composites alsoexhibit a higher temperature capability, lower creep, and increasedmodulus and hardness.

It is known that the intermetallics may deform by a number ofmechanisms, such as movement of dislocations, superdislocations,twinning, and the like. In the particular case of the intermetalliccompound TiAl, the lattice structure has an aspect ratio of 1.02, i.e.,the structure is nearly cubic. In this condition this structure maydeform by any of the aforementioned mechanisms.

The ambient temperature ductility of this material is determined by theease of the dislocation movement and it is therefore desirable tomaximize the number of operating slip systems to achieve the mostductile material. From a statistical point of view, it is known that areduction in the grain size of the intermetallic matrix will makeavailable a larger number of suitably oriented grains for dislocationactivity, an effective means of improving ductility.

A further means of increasing the number of available slip systems isthe attainment of a cubic structure, i.e., aspect ratio of 1.0. This maybe achieved by such means as alloying, radiation damage, and the like,such that the lattice is expanded in the shorter direction. The choiceof suitable alloying additions will be governed both by the size of theatomic nuclei and the electronic band structure of the alloyingelements. Examples where modification of the lattice parameter of anintermetallic by alloying is possible is the case of substitution of Alby Ti in Ni₃ Al to form Ni₃ (Al,Ti).

The combination of any of the above mechanisms for increasing thedeformation capabilities of intermetallics is considered a powerfulmeans of achieving enhanced ambient temperature ductility. When this iscombined with the dispersion strengthening and high-temperaturecapability afforded by the techniques disclosed herein, a unique seriesof composite materials is achievable which can satisfy both lowtemperature ductility concerns and also high temperature strengthrequirements.

For the purpose of illustrating the various reaction modes that may beused to form a second phase dispersion within an intermetallic matrix,detailed consideration will be given only to the specific case ofintermetallic-ceramic composites. In this discussion, it is understoodthat multiple dispersoids, intermetallic dispersoids, and/or multipleintermetallic matrices are also intended to be included. Methods ofpreparation of the intermetallic-ceramic composites of the presentinvention include the following: (A) coformation of the intermetallicand ceramic by inclusion of appropriate amounts of starting ingredients,in elemental form, so as to ultimately provide essentiallystoichiometric proportions of the constituents of the intermetallic andceramic, in a reaction vessel; (B) use of previously formedintermetallic material as a solid, followed by an in-situ precipitationof the ceramic material therein in a separate stage; (C) preparation ofa master concentrate of the ceramic phase dispersoid in one or moremetals or alloys that do not form intermetallics with themselves, whichmetals are convertible to form the desired intermetallic, followed bydilution of said concentrate in an intermetallic, or in further metalsor alloys with which the matrix material of the concentrate is reactiveto form intermetallics; (D) direct addition of one or more ceramicforming precursor materials in the presence of an intermetallic derivedsolvent material to a melt comprising an intermetallic derived solventand the complementary ceramic precursor(s) to generate an in-situceramic dispersoid forming reaction in the intermetallic derivedsolvent; and, (E) direct addition of one or more ceramic formingprecursors and an intermetallic precursor to a melt containing one ormore metals or alloys with which said intermetallic precursor isreactive to form intermetallics, and with which said ceramic formingprecursor is reactive to form the ceramic dispersoid.

As can be appreciated, a vast number of reaction sequences can beenvisioned as being within the scope of the present invention. Arepresentative number of such sequences are set forth in FIG. 1,although this schematic representation should not be taken as limitingthe scope of the present invention.

It is also recognized, moreover, that intermetallic-ceramic compositesmade by any of the methods described may be subjected to dilution or toaddition to melts of intermetallics or other matrix materials in whichthey are soluble, to achieve specific goals and purposes. For sake ofclarity, however, this procedure is not set forth in FIG. 1.

In the above-mentioned coformation process, (A), the starting materialsconstitute individual powders of each of the ceramic precursors and theintermetallic precursors. Thus, one may react a mixture of aluminum,nickel, titanium, and boron, to form a dispersion of titanium diboridein a nickel aluminide matrix. In this process, the basic reaction modemay be written as

    pM+qN+xA+yB→M.sub.p N.sub.q +A.sub.x B.sub.y,       (I)

wherein M and N represent ceramic precursor materials, A and B aremetallic elements constituting precursors of an intermetallic compound,and p,q,x and y closely approximate integers. In the specific examplegiven, the formation of titanium diboride in a matrix of nickelaluminide, M is titanium, N is boron, A is nickel, B is aluminum, p is1, q is 2, y is 1, and x may be 1 or 3, depending on which nickelaluminide, (NiAl or Ni₃ Al) is desired.

In this process, the intermetallic constituting the matrix material ofthe subject composite may be formed prior to, essentially simultaneouslywith, or after the in-situ precipitation of the ceramic dispersoid. Inpracticing this coformation process, it has been found that the natureand amount of the intermetallic and ceramic phase may be controlled byappropriate choice of the stoichiometry of the initial ingredientsadded. Thus, for example, a preponderance of Ni₃ Al or NiAl may beformed by correctly proportioning the relative amounts of the startingingredients, taking into account any of the ingredients consumed byceramic forming reactions, and losses due to absorption into the wallsof the reaction vessel, extraneous compound formation, volatilization,etc. In this regard, it should be noted that the ceramic formingconstituents must have a greater affinity for each other than for any ofthe intermetallic forming materials. Clearly, this only applies up tothe stoichiometric limit of the ceramic forming reaction, since beyondthis point, any excess of metallic ceramic forming constituent presentwill be available to form intermetallics. In many instances, theselection of the particular ceramic dispersoid and the particularintermetallic for the composite may dictate the sequence in which thecomponents are formed. For example, when preparing an aluminidecomposite where the ceramic and aluminide are formed in one operation,the aluminide may form at a lower reaction temperature and thus will beformed first. On the other hand, the ceramic may form at a lowertemperature and the aluminide form thereafter at a higher reactiontemperature. In the first instance, wherein the aluminide is formedfirst, the ceramic precursors are soluble in the molten aluminide, andprecipitate upon initiation of the ceramic-forming exothermic reaction.In the second instance, wherein the ceramic-forming reaction isinitiated at a lower temperature, the exothermic reaction liberatessufficient energy to raise the temperature of the reaction mass to thepoint at which the intermetallic precursors combine. Obviously, it ispossible for the temperature to rise sufficiently in some cases for bothreactions to occur essentially simultaneously.

In method (B) set forth above, one may use a previously preparedintermetallic, and thereafter add the necessary materials to form thedesired ceramic to a melt of the intermetallic. The basic reaction modecontemplated herein constitutes the following two-step procedure,wherein formula II represents the initial preparation of theintermetallic.

    xA+yB→A.sub.x B.sub.y                               (II)

    A.sub.x B.sub.y +pM+qN→M.sub.p N.sub.q in A.sub.x B.sub.y(III)

In a sense, this reaction mode may be considered a sub-set of Formula I,wherein the intermetallic forms prior to the ceramic dispersoid. In theformation of titanium carbide in nickel aluminide, A is nickel, B isaluminum, M is titanium, N is carbon, P is 1, q is 1, y is 1, and x maybe 1 or 3 dependent upon the aluminide desired.

In the ceramic master concentrate method (C), one forms a "master alloy"of ceramic dispersoid in a metal matrix, which matrix is subsequentlyconverted to the desired intermetallic. The first stage of thistwo-stage reaction made may be exemplified as:

    pM+qN+xA→M.sub.p N.sub.q +xA                        (IVa)

or

    pM+qN+yB→M.sub.p N.sub.q +yB                        (IVb)

The reaction products of Formula IV may then be processed in accordancewith Formula V, as follows:

    (M.sub.p N.sub.q +xA)+yB→M.sub.p N.sub.q +A.sub.x B.sub.y (Va)

or

    (M.sub.p N.sub.q +yB)+xA→M.sub.p N.sub.q +A.sub.x B.sub.y (Vb)

Alternatively, the "master alloy" of Formula IVa may be diluted in anintermetallic, so as to achieve a dispersion of ceramic in a matrixconsisting of mixed metal and intermetallic phases, or, in a matrixcomprising an intermetallic other than the diluent.

    (M.sub.p N.sub.q +xA)+A.sub.x' B.sub.y' →M.sub.p N.sub.q +A.sub.x' B.sub.y' +xA                                              (Vc)

    (M.sub.p N.sub.q +yB)+A.sub.x' B.sub.y' →M.sub.p N.sub.q +A.sub.x' B.sub.y' +yB                                              (Vd)

    (M.sub.p N.sub.q +xA)+A.sub.x' B.sub.y →M.sub.p N.sub.q +A.sub.x+x' B.sub.y                                                   (Ve)

    (M.sub.p N.sub.q +yB)+A.sub.x B.sub.y' →M.sub.p N.sub.q +A.sub.x B.sub.y+y'                                                (Vf)

This latter approach permits the preparation of such composites as TiB₂dispersed in Ni₃ Al, by the reaction sequence:

    Ti+2B+Ni→TiB.sub.2 +Ni

    2(TiB.sub.2 +Ni)+NiAl→2TiB.sub.2 +Ni.sub.3 Al

Obviously, this sequence mode is permissive of a great number ofvariants, enabling one to achieve a great variety of results utilizingsuch procedures. Thus it is possible to prepare "Master Concentrates",containing a ceramic phase, which may be utilized to introduce theceramic phase to a specified diluent in controlled fashion. Thus, forexample, one may prepare a master alloy of a high percentage of titaniumdiboride in an aluminide, and add metal or additional aluminide toachieve a composite having the desired composition.

The fourth reaction mode contemplated by this invention, method (D),envisions the direct addition of a ceramic forming precursor material,in the presence of a minor proportion of an intermetallic derivedsolvent, to a melt of the intermetallic phase and the complementaryceramic precursor.

    M.sub.p +(N.sub.q +A.sub.x B.sub.y)→M.sub.p N.sub.q +A.sub.x B.sub.y (VIa)

    N.sub.q +(M.sub.p +A.sub.x B.sub.y)→M.sub.p N.sub.q +A.sub.x B.sub.y (VIIb)

Another reaction method considered appropriate is the method identifiedhereinabove as method (E), wherein one or more ceramic precursors areadmixed or alloyed with an intermetallic precursor, and added to a meltcontaining the necessary ceramic and intermetallic precursors to reactto form the desired material:

    (pM+xA)+(qN+yB)→M.sub.p N.sub.q +A.sub.x yB         (VIIa)

    (pM+yB)+(qN+xA)→M.sub.p N.sub.q +A.sub.x B.sub.y    (VIIb)

    (pM+xA+x'A)+(qN+yB+y'B)→M.sub.p N.sub.q +A.sub.x B.sub.y +A.sub.x' B.sub.y'                                                  (VIIc)

    (pM+qN+xA+x'A)+yB+y'B→M.sub.p N.sub.q +A.sub.x B.sub.y +A.sub.x' B.sub.y'                                                  (VIId)

The method exemplified by Formulae VIIc and VIId illustrate thepreparation of an intermetallic-ceramic composite in which the matrix isa mixture of two differing intermetallic materials. The methodexemplified by Formulae VIIa and VIIb may be likened to the alloy-alloymethod set forth in patent application Ser. No. 662,928, filed Oct. 19,1984, of which this is a continuation-in-part, and which is incorporatedherein by reference. Similarly, the general concept of the solventassisted reaction of ceramic forming precursors is set forth in saidpatent application.

Varying amounts of ceramic may be incorporated into the compositematerial, depending upon the end use and the properties desired in theproduct. As previously noted, for dispersion strengthened materialshaving high modulus, one may utilize a preferred range of from about 10percent by volume to about 25 percent by volume. However, the ceramicvolume fraction may be varied considerably, so as to produce a compositewith the desired combination of properties, within the range of fromabout 0.5 percent by volume up to the point at which ductility issacrificed to an unacceptable extent. In contrast, cermet-likecomposites of up to about 95 percent or more by volume of ceramicmaterial in the aluminide containing matrix may be produced. Preferredranges for such materials will, of course, be dependent upon the desiredend use. It is possible to effectively tailor the composition to achievea spectrum of properties by controlling the proportions of the reactantand solvent materials.

Moreover, the various reaction modes may be initiated in differentphysical states. Thus, the elemental powders can initiate the process ina plasma arc or flame, or via diffusion of the reactive species throughthe liquid solvent, or, in cases where solid phase diffusion is rapid,in a solid state. In the case where two alloys are used, each containingan alloying element constituting a reactive component, the reaction canalso occur in the solid state, liquid state, gaseous or in a plasma arcor flame achieved, for example, by striking an arc between electrodes ofthe two alloys.

As was previously stated, the present invention provides for theformation of one or more finely dispersed precipitates in a matrix ofone or more intermetallic containing materials. It is important that thesecond phase precipitate material is not soluble in, or reactive with,the intermetallic derived solvent, while the constituents of the secondphase, individually, are at least sparingly soluble in the intermetallicderived solvent. Thus, the exothermic dispersion reaction mechanismdepends upon a certain amount of each second phase forming constituentdissolving and diffusing in the intermetallic derived solvent, and whilein solution (either liquid or solid state), reacting exothermically toform the insoluble precipitate rapidly as a very fine particulate. Theintermetallic derived solvent or solvent matrix material provides amedium in which the reactive elements may diffuse and combine. Once theinitial reaction has occurred, the heat released by the exothermicreaction causes additional diffusion of reactive components in thesolvent matrix material, and allows the reaction to proceed. During theinitiation and reaction extremely high temperatures may be achieved invery short periods of time. During this time frame, essentially all ofthe reactive constituents in the solvent, metal react to form theinsoluble second phase, which immediately precipitates.

The cool-down period following initiation of the reaction andconsumption of the reactive constituents may be important in achievingvery small dispersoid size, and limiting dispersoid growth. It is knownthat at high temperatures, it is possible for the second phase particlesto grow, e.g., by dissolution/precipitation or by agglomeration. Thisshould be avoided, because of the negative effect of large particlesizes on ductility. The cool-down or quenching of the reaction is, in asense, automatic, since once the second phase-forming constituents arecompletely reacted, there is no further energy released to maintain thehigh temperatures achieved. However, one may control the rate ofcool-down to a certain extent by controlling the size and/or compositionof the mass of material reacted. That is, large thermal masses absorbenergy, and cool down more slowly, thus permitting growth of largerparticles, such as may be desired for greater wear resistance, e.g., foruse in cutting tools. Thus, the temperature may be reduced from themaximum temperature attained to a temperature where grain growth isminimal. The incidence of particle growth will depend on the particularsecond phase being formed.

The reaction initiation temperature has generally been found to berelatively close to the melting temperature of the solvent matrixmaterial utilized in liquid state reactions. For example, in theproduction of titanium diboride in a titanium aluminide from theelemental powders, the reaction proceeds at a temperature around 660°C., near the melting point of aluminum. It should be noted that in theabsence of a solvent matrix material, the reaction of titanium and boronto form titanium diboride was not observed to proceed below atemperature of about 900° C., and frequently did not go to completion,and there was essentially no control over the particle size of the finalproduct. While one need not actually reach the melting temperatureinitially, one must achieve a temperature where substantial diffusion ofthe reactive species in the solvent matrix material can occur, eitherlocally or generally. It is also observed that, in some cases, as oneincreases the temperature it is possible for one of the startingconstituents to diffuse into a solvent matrix material, forming an alloytherewith having a lower melting temperature than the matrixintermetallic, and thus lowering the reaction initiation temperature.

It is also to be noted that with the basic process, one may cause thecomplex precipitation of a plurality of systems. Thus, it is possible toprecipitate complex phases, such as Ti(B₀.5 C₀.5), or alternatively, toprecipitate a mixture of titanium diboride and zirconium diboride in analuminide containing matrix.

It has been found that the powders need not be compacted prior tofiring, but doing so allows easier diffusion and thus initiation atlower temperatures. For instance, loose powder mixtures of aluminum,titanium and boron tend to react at slightly higher temperatures thanhighly compacted powders. This is due to localized melting, andincreased diffusion, which are possible when the powders are in closeproximity. In addition, compaction is advantageous because the compactmay be handled as a free-standing body requiring no containment vesselthat can generate contaminants in the system and can be destroyed by thethermal shock of the reaction. The impurity issue is especiallyimportant when an oxide refractory containment vessel such as zirconiais used in processing materials containing titanium which areparticularly susceptible to effects such as oxygen embrittlement.

Porosity of the final composite can be minimized by a vacuum degassingoperation prior to initiation of the reaction. The degree of vacummapplied and temperature of the degassing step is determined purely bythe kinetics of evaporation and diffusion of any absorbed moisture orother gasses. High vacuum and elevated temperatures aid the degassingoperation.

During heat up, the starting powders should be protected from extensiveoxidation due to exposure to the atmosphere, as this will restrict thediffusion of the components into the solvent matrix material, and thereaction should preferably be carried out under an inert gas to avoidoxidation at high temperatures. Further, this reduces the loss ofreactant or solvent species as volatile oxides or other oxidizedspecies.

The particle size of the powders utilized in the elemental powder modedoes not appear to be critical. It has been found, however, thatparticle size of the second phase reaction product can depend upon suchfactors as heat-up rate, reaction temperature, cool-down rate, andcrystallinity and composition of the starting materials. Appropriatepowder sizes may range from less than 5 microns to more than 200 micronsto facilitate mixing procedures and provide sufficiently small diffusiondistances. For economic reasons, one normally may utilize the largerparticle size powders. It has been found that the particle size of theprecipitated second phase in the matrix may vary from less than about0.01 microns to about 5 microns or larger, dependent upon such factorsas cited above.

As was mentioned previously, one embodiment of the subject process forforming intermetallic-second phase composite materials comprisesproviding a substantially molten or liquid mass containing the solventmatrix material, and then adding at least one of the constituents orelements of the desired second phase material with a minor proportion ofcompatible solvent material, to the molten mass. Upon the addition, thesolvent assisted in-situ precipitation reaction is initiated to form anddisperse finely divided particles of second phase material in thematrix. Thus, for example, the molten mass could contain one of theconstituents of the desired second phase material, e.g., as preformedalloy, and one or more constituents would subsequently be added.Alternatively, all of the constituents could be added to the molten masseither sequentially or simultaneously. One convenient mode of practicingthis process is to compact powders or chips of the unreactedconstituents of the second phase material with a minor proportion ofcompatible solvent material, and then add the compact of constituents tothe molten metal mass.

The constituent or constituents are added to the molten metal mass alongwith sufficient solvent metal to allow the reaction to easily proceed.With this procedure, the metal of the molten mass could be differentthan the added solvent metal and thus need not be a solvent for theconstituents. As above, this process allows for two or more metals inthe matrix phase which may then be converted to one or moreintermetallic phases by suitable heat treatment. In any such process,however, care must be taken that the diluent metal is not reactive withthe second phase or its constituents.

In this preferred addition process, it is generally preferable that theamount of each constituent added is such that essentially all of thesecond phase-forming constituents are consumed in the precipitationreaction, i.e., that essentially no unreacted second phase formingconstituent remains after the completion of the reaction. In mostinstances, this requirement can be met if stoichiometric quantities ofthe constituents are present in the final molten mass of matrix metaljust prior to solidification. In some cases, however, an excess of onecomponent beyond stoichiometric may be desirable, e.g., excess boron inthe preparation of TiB₂ in Ni₃ Al. Clearly, this stipulation is modifiedif one or more of the second phase-forming constituents is a metal thatis also reactive in forming the desired intermetallic matrix. Theamounts of the ingredients must then be proportioned according to thecombined stoichiometry of the required second phase precipitate and theintermetallic matrix phase.

In selecting the constituents and the matrix for the composite materialsproduced by the above-described addition process, it is important thatthe formed second phase material have low solubility in the molten mass.Otherwise, significant particle growth of the second phase material maybe experienced over extended periods of time at temperature. For mostapplications of the composite materials, the size of the second phaseparticles should be as small as possible, and thus particle growth isundesirable. When the solubility of the formed second phase material inthe molten mass is low, the molten mass with dispersed second phaseparticles can be maintained in the molten state for a considerableperiod of time without growth of the second phase particles.

An advantage of this embodiment is that, if the constituents are addedin a step-wise or incremental fashion, the bulk temperature of themolten mass will not change significantly during the course of theaddition, i.e., the large temperature spike associated with reactioninitiation and progress and thus potential particle growth of the secondphase particles due to elevated temperatures will be localized to thereaction zone, and rapidly quenched by the surrounding reaction mass,which acts essentially as an isothermal heat sink. Such an additionprocedure is also advisable from a safety standpoint to prevent therapid evolution of significant quantities of heat which could causemetal to be splattered or sprayed from the containment vessel. Anotheradvantage is that the agitation due to energy release and to temperaturegradients caused by the exothermic reaction of the constituents informing the second phase material occurring in the molten mass creates amixing effect and thus aids in dispersing the second phase materialthroughout the mass. In addition, by having the mass in the molten orliquid state upon addition of the constituents, the constituents arerapidly heated to reaction temperature, thus promoting the formation offine particles. A further important consideration of this process isthat since a molten mass of matrix metal is utilized, the matrix metalneed not be formed from powdered metal, a significant saving in materialpreparation costs.

As was also previously mentioned, one can prepare master concentrates ofthe subject composite materials and thereafter dilute the concentrate toyield the desired composite material. Generally, the concentrateformation comprises initially preparing a reactant mixture of secondphase forming constituents and then heating to produce the in-situreaction as described herein to form fine particles of second phasematerial dispersed in the matrix. Alternatively, the initialmetal-second phase composite can be formed by the previously describedprocess of adding the constituents directly to a molten mass of matrixmetal. The concentration or loading of second phase material isgenerally rather high, e.g., at least 10 volume percent, preferably 15percent, up to 80 or 90 volume percent or more of second phase materialin the resultant composite. Generally, concentrations below about 10volume percent are not economical for further dilution, andconcentrations in excess of about 90 volume percent are not advisable asthe reaction may become too violent and particle growth may beexperienced. After solidification, preferably the composite iscomminuted to a desired size, or alternatively, the composite can becast to any appropriate size.

The next step in the master concentrate process is dilution of theconcentrate by additional matrix material which can be the same ordifferent from the intermetallic forming material used in the in-situprecipitation reaction. In one embodiment, the diluting metal may beselected such that it forms one or more intermetallic phases with theoriginal matrix metal in which the precipitation reaction took place.Generally, this dilution may be accomplished by preparing a melt of theadditional matrix material and adding the composite to the melt, oralternatively, placing both the composite material and the additionalmatrix material in solid form in a vessel and then heating to atemperature such that the additional matrix material melts. Dispersionof the second phase material in the melt is facilitated by meltagitation generated by arc melting, mechanical stirring, inductionstirring, gas bubbling, ultrasonics, and the like. While in the moltenstate, various clean-up techniques such as the use of deoxidants,scavengers and the like can be employed to remove impurities such asoxygen from the matrix phase. Particularly advantageous embodimentsinclude the use of oxygen scavengers such as yttrium or erbium intitanium melts, thereby improving the scavenging of oxygen, and formingadditional finely dispersed oxides in the matrix. Once dispersion of thesecond phase material is complete, the melt may be solidified byconventional techniques such as chill casting to yield very low porositycomposites.

In preparing such master concentrates, degassing of the powders of theinitial reactant mixture may not be necessary, and in fact it may beadvantageous, from a processing standpoint, not to degas the powders,since a porous product tends to be formed which aids in the subsequentdilution by molten material. It even may be desirable in some instancesto incorporate a porosity enhancer such as a low boiling point metal,e.g., magnesium, in the initial reactant mixture, the enhancervolatilizing during the in-situ reaction, thereby increasing theporosity of the resultant composite. Use of hydrides may be particularlyadvantageous in this context since the hydrogen generated upondecomposition of the compound generates porosity and may be useful inreducing absorbed oxygen in the system.

The use of master concentrates, particularly those having high loadingsof second phase material, is advantageous since one can simply make onebatch of composite material and make a wide variety of differingcomposites having different dispersoid loadings. Another advantage isthat the additional matrix material used to form the melt need not be inpowder form, thereby saving considerably on raw material preparationcosts. Additionally, with the master concentrate procedure, it ispossible to form the second phase material in a matrix material whichis, for example, conducive to the formation of particles of a desiredsize, type, morphology, etc. and thereafter incorporate the particles ina compatible matrix material in which such particles cannot be producedby the in-situ precipitation reaction.

The following examples illustrate the precipitation of fine particles ofa dispersoid to produce a composite having an aluminide containingmatrix.

EXAMPLE 1

An intermetallic-ceramic composite containing about 35 weight percenttitanium diboride particles dispersed in a matrix of titanium aluminide(Al₃ Ti) is prepared as follows. A well-blended mixture of 202.5 gramsof aluminum, 239.5 grams of titanium and 55.7 grams of boron is madefrom powders of the respective elements and the mixture thenisostatically compacted with a pressure of about 35,000 psi. The formedcompact is heated in an inconel retort and a reaction initiated at about660° C., causing melting of the compact. Upon removal from the retort,the compact is subjected to X-ray analysis which indicates the presenceof TiB₂ and Al₃ Ti with only trace amounts of the initial elements. AnSEM analysis indicates that the titanium diboride particles aresubmicron and dispersed in a titanium aluminide matrix as is shown inFIG. 2. A sample where a portion of the Al₃ Ti matrix has been etchedaway is shown in FIG. 3 and gives further indication as to the fineparticle size and even dispersion of the titanium diboride. EDS analysisof the particles indicates that the particles are essentially puretitanium diboride.

EXAMPLE 2

An intermetallic-ceramic composite of titanium diboride particles in amatrix of titanium aluminide (Ti₃ Al) is prepared as follows. A mixtureof 67.5 grams of aluminum, 359.2 grams of titanium and 55.7 grams ofboron is thoroughly blended and the mixture then compacted and heated inthe manner of Example 1. The reaction temperature is observed to beabout 660° C. The resultant material upon solidification is a dispersionof fine particles of titanium diboride in a matrix of titanium aluminide(Ti₃ Al), as is shown in the photomicrograph of FIG. 4.

EXAMPLE 3

An intermetallic-ceramic composite containing 35 weight percent finetitanium diboride particles in a matrix of titanium aluminide (TiAl) isprepared as follows. A powdered mixture of about 117 grams powderedaluminum, about 328 grams titanium and about 56 grams of boron isprepared and mixed well to insure uniformity. The mixture is compactedand heated in the manner of Example 1 to yield a composite of finetitanium diboride particles in a matrix of TiAl as is shown in thephotomicrograph of FIG. 5. Analysis of the composite also reveals aminor amount of Ti₃ Al.

EXAMPLE 4

An intermetallic-ceramic composite of zirconium diboride particlesdispersed in titanium aluminide (Ti₃ Al) is prepared as follows. Amixture of 10.8 grams zirconium, 2.5 grams boron, 76.6 grams titaniumand 13.8 grams aluminum is thoroughly blended and then processed in amanner similar to that of Example 1 to yield the composite.

EXAMPLE 5

An intermetallic ceramic composite of zirconium diboride particlesdispersed in a matrix of nickel aluminide (Ni₃ Al) is prepared. Twomixtures, each containing 97 grams of zirconium, 23 grams of boron, 243grams of nickel and 37.2 grams of aluminum are blended thoroughly. Onemixture is heated to reaction initiation temperature in a resistanceheated furnace, and the other heated to reaction initiation temperatureby induction heating. The resultant composites each contain a smallamount of unreacted nickel. When subjected to fracturing forces, thesecomposites have a fracture surface which exhibits microvoid coalescence,which tends to indicate that the mode of fracture was a ductile one,consistent with the fine grain size of the aluminide matrix. Thefracture surface is shown in the photomicrograph of FIG. 6.

EXAMPLE 6

A mixture of nickel, aluminum, titanium, and boron in the stoichiometricproportions for the formation of nickel aluminide (Ni₃ Al) and titaniumdiboride (TiB₂), i.e., 10 percent by weight aluminum, 62 percent byweight nickel, 19 percent by weight titanium, and 9 percent by weightboron, is compacted to 40,000 pounds per square inch, and then heated ina furnace. Upon reaching 620° C., a rapid exotherm is noted, whichsubsequent analysis by X-ray diffraction and scanning electronmicroscopy identifies as resulting from the formation of titaniumdiboride particles in a nickel aluminide matrix. It is evident from thisexperiment that a ceramic phase, e.g., titanium diboride, could bedirectly preciitated in an intermetallic phase, e.g., nickel aluminide,provided the affinity of the ceramic-forming species for each other isgreater than either has for the two elements making up the intermetallicmatrix.

EXAMPLE 7

An intermetallic-ceramic composite of titanium diboride particlesdispersed in a matrix of nickel aluminide (Ni₃ Al) is prepared asfollows. A mixture of 103.5 grams titanium, 46.5 grams boron, 302.5grams nickel, and 47.5 grams aluminum is blended and then isostaticallypressed. About 100 grams of the pressed compact is then reacted in aretort. From a temperature probe placed adjacent to, but not touchingthe compact, the reaction apparently occurs at about 807° C. and thetemperature during reaction peaks at about 1050° C. X-ray diffraction ofthe resultant composite indicates the presence of TiB₂, Ni₃ Al, andresidual Ni.

EXAMPLE 8

In a series of experiments, the formation of each of the dispersoidshafnium carbide, zirconium carbide, titanium carbide, titanium boride,titanium diboride, and vanadium diboride in the matrices of titaniumaluminide (Ti₃ Al) and nickel aluminide (Ni₃ Al) is investigated. Inpreparing the various composites, the constituents forming the ceramicdispersoid and the components forming the aluminide containing matrixare reacted at the same time. The constituents and components in thereacting mixture are combined in such proportions so as to yield anintermetallic matrix composite containing about 40 weight percentceramic dispersoid. The reactions for each composite are conductedtwice, one of the reactions being conducted under an argon atmosphereand the other under vacuum. Induction heating is used to initiate eachreaction, and at the first indication of a reaction, power to theinduction heating unit is terminated so that the composite may cool asquickly as possible.

Upon completion of the reaction, each of the formed ceramic-aluminidecomposites is examined by X-ray diffraction analysis to determine itscomposition. In addition, a small amount of the matrix is dissolved inacid and the ceramic particles are observed for particle size by ascanning electron microscope and also examined by X-ray diffraction todetermine the particle composition.

The results of these observations are set forth in the following Table.

                                      TABLE I                                     __________________________________________________________________________              Reacted Under Argon                                                                             Reacted Under Vacuum                              Desired Compound                                                                        Major  Minor Particle                                                                           Major  Minor Particle                             Dispersoid/Matrix                                                                       Component                                                                            Component                                                                           Size Component                                                                            Component                                                                           Size (micron)                                                                        Comment                       __________________________________________________________________________    Hfc/Ni.sub.3 Al                                                                         HfC, Ni.sub.3 Al,                                                                    Hf, Al.sub.2 Hf                                                                     0.1-2                                                                              HfC, Ni.sub.3 Al                                                                     Hf    0.1-1.3                                                                              Amount of Al.sub.2 Hf                   Ni                Ni                  was very small                HfC/Ti.sub.3 Al                                                                         HfC, TiC                                                                             Ti.sub.3 AlC                                                                        0.1-1.2                                                                            HfC, TiC                                                                             Ti.sub.3 AlC                                                                        0.5                                            Ti.sub.3 Al, Ti   Ti.sub.3 Al, AlTi.sub.2                           VB.sub.2 /Ni.sub.3 Al                                                                   VB.sub.2, Ni, Al,                                                                      --  0.1-1.3                                                                            VB.sub.2, Ni.sub.3 Al,                                                               V     1-2                                            V.sub.3 B.sub.4   V.sub.3 B.sub.4, Ni, Al                           VB.sub.2 /Ti.sub.3 Al                                                                   Ti.sub.3 Al,                                                                         V, TiB.sub.2                                                                        none Ti.sub.3 Al,                                                                         TiB.sub.2,                                                                          none                                           V.sub.3 V.sub.4   Al.sub.11 V                                                                          Al.sub.6 V                                           Al.sub.11 V,                                                                  Al.sub.6 V                                                          TiC/Ni.sub.3 Al                                                                         TiC, Ni.sub.3 Al                                                                       --  0.1-2                                                                              TiC, Ni.sub.3 Al,                                                                    --    0.1-1.5                                        Ni                Ni                                                TiC/Ti.sub.3 Al                                                                         TiC, Ti.sub.3 AlC                                                                    Ti.sub.3 Al                                                                         0.1-1.5                                                                            TiC, Ti.sub.3 AlC                                                                      --  0.2-2                                ZrC/Ni.sub.3 Al                                                                         ZrC, Ni.sub.3 Al,                                                                    Ni.sub.5 Zr                                                                         1    ZrC, Zr, Ni,                                                                           --  0.5    Intermetallics of Ni and                                                      Zr are                                  Ni.sub.7 Zr.sub.2 Ni.sub.3 Al         probably unstable phases      ZrC/Ti.sub.3 Al                                                                         TiC, Ti.sub.3 Al,                                                                    AlZr.sub.3                                                                          none Ti.sub.3 AlC, TiC,                                                                     --  none   Did not form ZrC                                                              dispersoid,                             Ti.sub.3 AlC      Al.sub.2 Zr         due to greater stability                                                      of TiC,                                                                       Ti.sub.3 AlC                  TiB/Ti.sub.3 Al                                                                         Ti.sub.3 Al                                                                          TiB   Ti.sub.2 Al                                                                          --     --  --                                   (10 vol %)                                                                    TiB/Ti.sub.3 Al                                                                         Ti.sub.3 Al                                                                          TiB   Ti.sub.2 Al                                                                          --     --  --                                   (30 vol %)                                                                    TiB and TiC/                                                                            TiC      --  --     --     --  --     Using B.sub.4 C as                                                            reactant                      Ti.sub.3 Al                                                                             TiB                                                                           Ti.sub.3 Al                                                         TiB.sub.2 /TiAl                                                                         TiAl   TiB.sub.2                                                                           --     --     --  --                                   TiC/TiAl  TiC    Ti.sub.3 Al                                                                         --     --     --  --                                             TiAl   Ti.sub.2 Al                                                  __________________________________________________________________________

EXAMPLE 9

An intermetallic-ceramic composite having mixed ceramic dispersoids isprepared by mixing 11.0 grams of Al₄ C₃, 33.8 grams of tantalum, and135.2 grams of niobium, and heating in a graphite induction furnace.Analysis of the recovered product reveals the presence of both TaC andTa₂ C in a matrix of Nb₃ Al.

EXAMPLE 10

An intermetallic-ceramic composite comprising a ceramic dispersoid in amixed intermetallic matrix is prepared by mixing 43.6 grams of titanium,123.6 grams of tantalum, and 32.8 grams of Al₄ C, compacting, andreacting on a water cooled copper holder in an induction furnace underflowing argon. Upon recovery of the reaction product, X-ray analysisshows the presence of TiC and a mixed matrix of TiAl, TaAl₃, and TaAl₂.

EXAMPLE 11

An intermetallic-ceramic composite of titanium diboride in a matrix oftitanium aluminide (Al₃ Ti) is prepared by the master concentrate route.A solidified melt comprising 30 weight percent titanium diboride in atitanium aluminide matrix is comminuted to particles having an averagesize of about 1 millimeter, and then a melt of about 860 grams oftitanium aluminide is prepared under a protective inert atmosphere. Theparticles are then added to the melt and held at that temperature for asufficient period of time to insure complete melting of theintermetallic phase of the particles, and a uniform distribution of thetitanium diboride. The melt is then solidified to yield a composite of15 weight percent titanium diboride dispersed in a titanium aluminidematrix.

An intermetallic-ceramic composite is prepared by the direct additionroute, by mixing 65.5 grams of titanium, 10 grams of boron, and 24.2grams of aluminum, compacting, and incrementally adding to a molten poolof TiAl under inert atmosphere. On addition of the compact to the moltenpool, a reaction occurs resulting in the formation of fine, evenlydispersed TiB₂ particles. Upon completion of the addition, the mixtureis cast and recovered as a dispersion of TiB₂ in a matrix of TiAl.

EXAMPLE 12

A mixture of appropriate amounts of titanium, silicon, and copperpowders to form 60 volume percent TiSi₂ is compacted and subsequentlyheated under an inert atmosphere to initiate a reaction and theprecipitation of TiSi₂ in a copper matrix. The resultant concentrate isadded to molten gold to produce a composite having TiSi₂ particles in aCu₃ Au matrix.

In some cases where the intermetallic matrix phase formation does not goto completion, it may be desirable to homogenize the final product inthe solid state to complete the conversion of the intermetallicprecursors to the intermetallic. For example, in the formation of TiB₂in TiAl from the elemental powders, it is found that on recovery of thefinal product, in addition to TiAl, small amounts of Ti₃ Al and Ti₂ Alare present. Subsequently, homogenization of this composite at 1000° C.for 3 hours essentially eliminates all intermetallics except TiAl.

It should be noted that the process disclosed herein for making secondphase containing composites with an intermetallic containing matrix hasa number of advantages over powder metallurgical methods taught by theprior art for preparing intermetallic materials. For example, thepresent process circumvents the need for submicron, unagglomeratedsecond phase or intermetallic starting materials, which materials arenot normally commercially available, and which are often pyrophoric. Thepresent process also eliminates the technical problems of uniformlydispersing a second phase in an intermetallic, and avoids the problem ofoxides at the intermetallic/intermetallic or second phase/intermetallicinterface during processing. Further, the process yields anintermetallic-second phase composite with a second phase precipitatedin-situ therein, having one or more of the following properties:superior hardness, ductility, and modulus qualities superior tocurrently available intermetallic containing materials. The compositesalso have improved high temperature stability, since the second phase isselected such that it is not reactive with the matrix, and thus thecomposites can be welded while maintaining uniformly dispersed discretefine particles, and the resultant weldment possesses superior corrosionresistance when compared to the welded metal matrix composites presentlyavailable.

It is understood that the above description of the present invention issusceptible to considerable modification change, and adaptation by thoseskilled in the art, and such modifications, changes, and adaptations areintended to be considered to be within the scope of the presentinvention, which is set forth by the appended claims.

We claim:
 1. A method for the preparation of intermetallic-second phasecomposite materials, said method comprising: contacting reactiveprecursors of a second phase material and a solvent matrix materialcomprising an intermetallic or precursors thereof at a temperaturesufficient to permit diffusion of said second phase precursors into saidsolvent matrix materials and to initiate the exothermic reaction of saidsecond phase precursors; permitting the temperature to rise as a resultof said reaction to enable solvent assisted formation of the secondphase in said solvent matrix material; and recovering a compositecomprising an intermetallic containing matrix having particles of saidsecond phase dispersed therein.
 2. A method as set forth in claim 1,wherein said reactive precursors of said intermetallic are aluminum,nickel, copper, titanium, cobalt, iron, platinum, gold, silver, niobium,tantalum, boron, lead, zinc, molybdenum, yttrium, hafnium, tin,tungsten, lithium, magnesium, beryllium, thorium, silicon, chromium,vanadium, zirconium, manganese, scandium, lanthanum, rare earthelements, or alloys thereof.
 3. A method as set forth in claim 2,wherein said second phase comprises an intermetallic material other thanthe matrix intermetallic.
 4. A method as set forth in claim 1, whereinsaid second phase comprises a ceramic.
 5. A method as set forth in claim1, wherein the reactive second phase precursors and intermetallicprecursors are each provided as individual elements.
 6. A method as setforth in claim 1, wherein said reactive second phase precursors andintermetallic precursors are provided as alloys.
 7. A method as setforth in claim 1, wherein at least one reactive precursor of said secondphase material is a transition metal of the third to sixth group of thePeriodic Table.
 8. A method as set forth in claim 1, wherein at leastone second phase precursor is aluminum, titanium, silicon, boron,carbon, sulfur, molybdenum, tungsten, vanadium, zirconium, niobium,cobalt, nitrogen, oxygen, nickel, iron, magnesium, beryllium, manganese,zinc, lithium, yttrium, hafnium, tantalum, chromium, thorium, arefractory metal, a rare earth metal, or a reactive compound thereof. 9.A method as set forth in claim 8, wherein at least one intermetallicprecursor is aluminum, nickel, titanium, cobalt, iron, or a refractorymetal.
 10. A method as set forth in claim 9, wherein said second phaseprecursors are titanium, zirconium, hafnium, boron, silicon, oxygen,nitrogen, or carbon.
 11. A method as set forth in claim 10, wherein theintermetallic is Ti₃ Al.
 12. A method as set forth in claim 11, whereinthe second phase precursors are titanium and boron.
 13. A method as setforth in claim 11, wherein the second phase precursors are titanium andcarbon.
 14. A method as set forth in claim 10, wherein the intermetallicis TiAl.
 15. A method as set forth in claim 14, wherein the second phaseprecursors are titanium and boron.
 16. A method as set forth in claim14, wherein the second phase precursors are titanium and carbon.
 17. Amethod as set forth in claim 10, wherein the intermetallic is TiAl₃. 18.A method as set forth in claim 17, wherein the second phase precursorsare titanium and boron.
 19. A method as set forth in claim 17, whereinthe second phase precursors are titanium and carbon.
 20. A method as setforth in claim 10, wherein the intermetallic is NiAl.
 21. A method asset forth in claim 10, wherein the intermetallic is Ni₃ Al.
 22. A methodas set forth in claim 21, wherein the second phase precursors aretitanium and boron.
 23. A method as set forth in claim 21, wherein thesecond phase precursors are zirconium and carbon.
 24. A method as setforth in claim 21, wherein the second phase precursors are zirconium andboron.
 25. A method as set forth in claim 1, wherein plural second phasematerials are produced.
 26. A method as set forth in claim 25, whereinthe second phase precursors are compounds selected from boron carbide,boron nitride, and boron oxide.
 27. A method as set forth in claim 25,wherein one of said second phase materials is an oxide or a nitride. 28.A method as set forth in claim 27, wherein said oxide is an oxide ofyttrium, cerium, erbium, or a rare earth element.
 29. A method as setforth in claim 1, wherein said intermetallic containing matrix comprisesa mixture of intermetallic materials.
 30. A method as set forth in claim1, wherein said reactive second phase precursors and said solvent matrixmaterial are added to molten intermetallic material.
 31. A method as setforth in claim 30, wherein said molten intermetallic material is otherthan the intermetallic of the solvent matrix material.
 32. A method asset forth in claim 1, wherein said reactive second phase precursors andsaid solvent matrix material are added to molten intermetallicprecursor.
 33. A method as set forth in claim 1, wherein at least one ofsaid intermetallic precursors is a compound.
 34. A method as set forthin claim 1, wherein at least one second phase precursor is a hydride.35. A method for forming an intermetallic material having a finelydivided second phase material precipitated in-situ therein, said methodcomprising the steps of: contacting reactive precursors of second phasematerial, in the presence of reactive precursors of intermetallicmaterial in which said second phase material precursors are more solublethan said second phase material; raising the temperature of saidreactive precursors to a temperature at which sufficient diffusion ofsaid precursors of said second phase materials into at least one of saidreactive precursors of said intermetallic material occurs to cause theinitiation of a solvent assisted exothermic reaction of the precursorsof said second phase, thereby forming second phase particles in-situ;forming the intermetallic material from the precursors thereof; andrecovering a composite comprising an intermetallic containing matrixwith said second phase particles dispersed therein.
 36. A method as setforth in claim 35, wherein at least one intermetallic precursor isaluminum, nickel, titanium, cobalt, iron, or a refractory metal.
 37. Amethod as set forth in claim 36, wherein at least one of said secondphase precursors is selected from titanium, zirconium, hafnium, boron,silicon, oxygen, nitrogen, and carbon.
 38. A method as set forth inclaim 35, wherein said reactive precursors are added to moltenintermetallic material.
 39. A method as set forth in claim 38, whereinsaid molten intermetallic material is other than the intermetallicformed by the reactive intermetallic precursors.
 40. A method as setforth in claim 35, wherein said reactive precursors are added to moltenintermetallic precursor.
 41. A method for making intermetallic-secondphase composites, said method comprising: contacting second phaseforming constituents and an intermetallic material in which said secondphase forming constituents are more soluble than the second phase;heating to a temperature at which sufficient diffusion of said reactivesecond phase forming constituents into an intermetallic derived solventmaterial occurs to cause an exothermic solvent assisted reaction of saidconstituents, thereby precipitating second phase particles in-situ insaid intermetallic derived solvent; and recovering a material comprisingfinely divided second phase particles in an intermetallic containingmatrix.
 42. A method as set forth in claim 41, wherein the temperatureis at least about the melting temperature of said intermetallicmaterial.
 43. A method as set forth in claim 41, wherein at least oneintermetallic precursor is aluminum, nickel, titanium, cobalt, iron, ora refractory metal.
 44. A method as set forth in claim 43, wherein atleast one of said second phase forming constituents is selected fromtitanium, zirconium, hafnium, boron, silicon, oxygen, nitrogen, andcarbon.
 45. A method as set forth in claim 41, wherein said constituentsand said intermetallic material are added to molten intermetallic.
 46. Amethod as set forth in claim 41, wherein said constituents and saidintermetallic material are added to molten intermetallic precursor. 47.A method as set forth in claim 41, wherein said second phase formingconstituents and said intermetallic material are provided as premixedand compacted powders.
 48. A method for making intermetallic-secondphase composites, said method comprising: contacting a reaction mixtureconsisting of second phase forming constituents and a first precursor ofan intermetallic material, said second phase forming constituents beingmore soluble in said intermetallic precusor than said second phase;heating said reaction mixture to a temperature at which sufficientdiffusion of said second phase forming constituents into said firstintermetallic precursor occurs to initiate an exothermic solventassisted reaction of said constituents and to form a second phaseprecipitate in-situ in said first intermetallic precursor; adding saidsecond phase precipitate in said first intermetallic precursor to amolten mass of a second intermetallic precursor to thereby form anintermetallic having second phase particles therein; and recovering amaterial comprising finely divided second phase particles in anintermetallic containing matrix.
 49. A method as set forth in claim 48,wherein the temperature is at least about the melting temperature ofsaid first intermetallic precursor.
 50. A method as set forth in claim48, wherein said second phase precipitate in said first intermetallicprecursor is added to molten intermetallic.
 51. A method as set forth inclaim 48, wherein a solid mass of said second intermetallic precursor ismixed with said second phase precipitate in said first intermetallicprecursor, and heated to form the intermetallic.
 52. A method as setforth in claim 48, wherein said reaction mixture is added to a moltenmass of said first intermetallic precursor to initiate said reaction.53. A method as set forth in claim 52, wherein a solid mass of saidsecond intermetallic precursor is mixed with said second phaseprecipitate in said first intermetallic precursor, and heated to formthe intermetallic.
 54. A method as set forth in claim 52, wherein saidsecond phase precipitate in said first intermetallic precursor is addedto molten intermetallic.
 55. A method as set forth in claim 48, whereinsaid reaction mixture is added directly to an intermetallic derivedliquid.
 56. A method as set forth in claim 48, wherein said reactionmixture is added directly to a molten mass of a second intermetallicprecursor to initiate said reaction.
 57. A method as set forth in claim48, wherein at least one intermetallic precursor is aluminum, nickel,titanium, cobalt, iron, or a refractory metal.
 58. A method as set forthin claim 57, wherein at least one of said second phase precursors isselected from titanium, zirconium, hafnium, boron, silicon, oxygen,nitrogen, and carbon.
 59. A method for making intermetallic-second phasecomposites, said method comprising; adding one or more second phaseprecursors to a molten reaction mixture comprising an intermetallicderived solvent and at least one complementary second phase precursor;forming said second phase by an exothermic solvent assisted reaction insaid molten reaction mixture; and recovering a material comprising adispersion of second phase particles in an intermetallic containingmatrix.
 60. A method as set forth in claim 59, wherein at least oneintermetallic precursor is aluminum, nickel, titanium, cobalt, iron, ora refractory metal.
 61. A method as set forth in claim 60, wherein saidsecond phase precursors are titanium, zirconium, hafnium, boron,silicon, oxygen, nitrogen, or carbon.
 62. A method for makingintermetallic-second phase composites, said method comprising: addingone or more second phase precursors and an intermetallic precursor to amelt of one or more intermetallic precursors and at least onecomplementary second phase precursor; permitting an exothermic solventassisted reaction of the respective precursors; and recovering adispersion of second phase particles in an intermetallic containingmatrix.
 63. A method as set forth in claim 62, wherein at least oneintermetallic percursor is aluminum, nickel, titanium, cobalt, iron, ora refractory metal.
 64. A method as set forth in claim 63, wherein saidsecond phase precursors are titanium, zirconium, hafnium, boron,silicon, oxygen, nitrogen, or carbon.